Aeroengine Safety

Figure "Flaws in forged rotor disks" (Ref. 15.2-15): This diagram depicts the most common flaw types in forged parts with their effects on the part properties. It also includes any relevant references to sections that deal with the specific problems in greater detail.

“Grain structure”: This term is confusing. In materials of “engine quality”, it merely refers to structural orientation. This is primarily determined by a favorable grain boundary pattern and the location of carbides in the grain.
Unlike materials such as construction steel, which have a considerable proportion of material-typical inhomogeneities (not flaws!), the high-strength materials used in engine construction are especially pure. However, even here there is an influence of grain boundary patterns (e.g. under creep loads, plastic deformability) and any inhomogeneities that affect strength (weak points). One attempts to minimize negative influences, especially those that affect dynamic fatigue and creep life, through favorable grain structure orientation (texture), i.e. parallel to the highest operating loads. This orientation is achieved with the aid of plastic deformation during the forging process (Fig. "Formation of forging flaws"). It also corrects flaws such as unallowable segregations. However, the altered location of these flaws can make them more difficult to detect with ultrasonic testing (Ills. 15.2-21 and 17.3-15). Therefore, non-destructive testing must be optimized accordingly. It is not always possible to guarantee sufficient forged re-forming, especially in thick cross-sections such as the hub area of a disk. This promotes dangerous flaws precisely in this especially highly stressed part zone. This includes all flaws from melting (Fig. "White spots in Ni alloys") and casting porosity that hasn`t been completely forged. An additional effect is the smashing of larger inclusions into many smaller, less dangerous small ones (Fig. "Safety of parts by manufacturing"). Therefore, unfavorable grain structure and/or insufficient forging deformation increase the risk of life-reducing flaws. The cyclical life span (LCF) is especially affected.

Crack initiation: “Forging cracks” usually occur near the surface. They are caused by friction forces between tools (forging hammer, forging die, extrusion die) and the part being forged. Internal forging cracks that are not oxidized will usually disappear by being welded shut during the forging process. Common crack locations are in the hub area. Many forging cracks in the surface zone of finished parts can be avoided by ensuring that there is a sufficient oversize in the raw part. The remainder should be sufficiently safely detectable through serially implementable non-destructive testing methods (e.g. ultrasonic, eddy current, and penetrant) that have been optimized for the specific parts. The trend toward ever higher utilization of material strength makes the limits of detection for these processes even tighter (Fig. "Testing of high strength material").
Heat treatment cracks (Ills. 15.1-19 and 15.2-13) occur as a result of overly high thermal stress and insufficient strength in the heating or cooling phases. If the cracks oxidize or if the deformation is insufficient for them to reseal during the forging process, these cracks will not be “healed”.

Deformation-dependant strength losses occur when sufficient deformation does not occur during forging. This is especially common in thick hub cross-sections that have a high resistance to solid forging. This means that the required static and dynamic minimum strength values may not be met in local areas. The reason for this is less than optimal development of the structure (Fig. "Specific part structure"). Especially affected materials are ones in which the strength must be ensured through a strictly specified thermomechanical reforming process. It primarily affects LCF strength and creep resistance.

Strength losses through unfavorable temperature control: the larger the part cross section, the slower temperature changes will occur (Fig. "Forging process caused characteristics"). This is especially true of titanium alloys with their low thermal conductivity. However, even in raw disk parts made from Ni alloys with large masses in the hub area, as are typical for turbine disks of small gas turbines (middle diagram), there is a danger that the structure will not be optimal (Fig. "Solution annealing and gamma phase").

Residual stresses are created in parts during the forging process through deformation and/or restricted thermal strain during subsequent heat treatment. If tension residual stresses occur in a part area that is subject to high operating loads, they can unallowably shorten the life span of the part (Fig. "Residual process stresses in titanium rotor disk"). The greater the thermal strength of the part, the greater the residual stresses that cannot be broken down through stress-relieved annealing, and these can combine with the operating stresses in the part (Fig. "Residual process stresses in titanium rotor disk").

Thermomechanical Processing (TMP) can create structural zones that appear light during etching and are referred to as “pseudo white spots” in the technical literature (Ref. 15.1-4). However, these are not inhomogeneities of alloy components that present a cause for concern, but are rather coarse grain zones (detail in bottom frame) that are surrounded by very fine grain. These coarse grain zones are created when the annealing temperature does not completely reach the level required for complete recrystallization. The difference may only be a few °C.
This phenomenon is especially known from the material IN 718, which has a broad temperature range between 995°C, where crystallization begins, and 1030°C, at which point the entire structure has been reached.
The recrystalization model (bottom frame) uses a structured angled plane to show how the individual structural zones recrystallize when they reach the necessary energy state (temperature).

Figure "Solution annealing and gamma phase" (Ref. 15.2-5): the development of the g' phase, which is responsible for thermal strength, depends largely on its solution and separation during cooling. In this process, the temperature levels during heat treatment are decisive. Especially important are the heating and cooling rates in the entire part, as well as strict adherence to the solution annealing temperature. Even small deviations from the solution annealing temperature of only 2% can, for example, change the important g' component by over 20% in the depicted high-strength Ni forged alloy, which has a decisive influence on thermal strength.

Figure "Disk cracking during heat treatment": During a testing rig run in the development phase of a military engine with multiple shafts, a serious imbalance occurred. Disassembly revealed a gaping circumferential crack about 20 centimeters long on a turbine disk at the flange connection to the disk (right diagram). Unlike the disk surface, the open crack surface was heavily oxidized. A metallographic cut confirmed the findings (top detail). This oxidation could not be explained with the known operating conditions. A fracture surface analysis did not show any dynamic fatigue crack signs, except for the ends of the crack. The crack surface, as far as it could be analyzed, revealed the typical “doughy” characteristics of a thermal crack (Ill.15.1-8). Further research revealed that the crack occurred during heat treatment of the part in its raw part contour, and also occurred in several additional raw parts (left diagram). The raw part contour had a notch and abrupt cross-section jump to a rapidly heating thin ring in the area where the cracking later occurred. This allowed dangerously high thermal stress to develop during heating and cooling. It is not known why the part was finished and installed in the engine despite the large crack length (Fig. "Failed detection of large cracks"). The decisive factor may have been the “horizon of expectation” of the crack testers, which was geared to much smaller cracks, combined with the oxides filling the crack.

Figure "Titanium flaw types" (Ref. 15.2-30): The Sioux City flight accident (Fig. "Problems of titanium alloy rotor raw parts") was traced back to a material flaw in the fan disk of one of the three engines. The material was a titanium alloy. A review team made up of specialists was tasked with evaluation of the entire technology, including raw part production, construction, and design configuration. They were also to determine possible suggestions for improvements. In this context, 22 relevant damages and incidents that occurred before 1990 were compiled and analyzed.
The causal material flaws could be classified into four categories corresponding to their microstructure, chemical analysis, and the physical characteristics that influenced the fracture surfaces:

Type I/Category 1: This flaw is relatively common and is related to process controls and quality problems. It occurs due to “burned” (reaction with air) titanium sponge from the melting process. The typical result is three defined concentric zones (top left diagram) Zone 1 in the inside has a sponge-like structure with large pores with an average diameter of 0.75 mm. Zone 2, which only occurs in this flaw type, consists of very hard and probably also brittle nitrogen-stabilized a structure. Zone 3 is the same in all four categories. It is an ellipsoid of a structure, with its orientation corresponding to the forging deformation, and surrounded by a-accumulations or a-plates.

Type I/Category 2: This flaw type is also relatively common and is related to process controls and quality problems. Zone 1 can contain pores, surrounded by a zone 3 made of nitrogen-stabilized a structure. The hardness is R_C 55-70, which is clearly higher than the normal R_C 35-40. The pores have an average diameter of 0.1 mm, making them considerably smaller than in Category 1. Pores of such a small size evidently do not explain the crack growth in damage cases. Apparently, the entire flaw zone has a weakening effect. Here, as well, zone 3 consists of nitrogen-stabilized a structure.

Type II/Category 3: Is considerably rarer than categories 1 and 2. It consists of a zone 1 with very small micropores with diameters of between 0.25-0.005 mm. Here, as well, zone 3 is made up of a structure, although in this case is it aliminum-stabilized. The hardness is R_C 35-40, thus corresponding to the normal values of the base material.

Type II/Category 4: These flaws have been the least common as damage causes. They consist merely of a single zone 3. These segregations consist of pure titanium and have a correspondingly low hardness/strength (R_C 12). Porosity has not been found. Flaws of this type were found before the introduction of “hot topping”.

Crack growth in the cracking zone in the investigated damage cases showed the following typical characteristics:

• crack initiation at pores
• smooth cracked surfaces (cleavage cracks) directly around the pores correspond to brittle micro-cracking (Fig. "Pores causing cracks in titanium alloys").
• inter-crystalline, facet-like cracking along plate-shaped a structures.
• LCF cracking and ductile splitting between the facets.

Depending on the flaw type, the fracture surface characteristics had varying degrees of pronouncement. In the case of category 1, porosity evidently had the dominant influence on crack growth. In categories 2 and 3, a similar structural influence could be discerned. An important realization was the fact that the pore size does not necessarily correspond to the effective crack initiation flaw size. For this reason, in the case of risk assessments, the assumption of a larger effective pore size seems to be necessary (Fig. "Flaw size and risk assessment").

Figure "Flaw size and risk assessment"

Figure "Titanium rotor disk flaw distribution" (Ref. 15.2-30): The analysis of flaws in rotor parts made from a titanium alloy in the time period before 1990 resulted in the following conclusion: the frequency of flaws from the melting process (Fig. "White spots in Ni alloys") and the forging process (Ills. 15.3-12 and 15.3-13) is not equal throughout the entire part volume. One can see that dangerous flaws occurred, or at least were discovered, especially in the hub area. This type of flaw distribution is also valid for Ni-based alloys and high-alloy steels produced in similar processes. Flaws in the hub area are especially relevant to safety. This area is typically subject to very powerful cyclical loads. In addition, in the thick hub cross-sections, it is more difficult to realize the plastic deformations during forging that are required to achieve the necessary strength properties (Ills. 15.1-14 and 15.2-11). This deformation is a prerequisite for the smashing of flaws (impurities, also see Fig. "Safety of parts by manufacturing") and/or giving them a favorable orientation relative to the main load direction (Fig. "Formation of forging flaws"). If the deformation is not sufficient, larger flaws in more effective positions are more likely.
About 80% of the burst rotors had crack initiation weak points below the surface. The LCF cracks grew towards the surface. This meant that dangerously large cracks could only be detected at an advanced stage through penetrant testing. Of course, it was a prerequisite that this type of testing occurred during this growth phase of the cracks, such as during overhauls.
Generally, it can be said that the probability of material flaws increases along with the volume of the parts, i.e with the thickness of the cross-sections.

Figure "Flaw causing titanium rotor fracture" (Ref. 15.2-18, Example 15.2-3): The one-piece compressor rotor (spool) made from the high-strength titanium alloy Ti6242 failed between the third and ninth stages (middle right diagram) due to LCF (bottom left diagram). The crack was traced back to a decrease in dynamic strength in the area of an oxygen accumulation (segregation, bottom right diagram) with an increased proportion of a-structure. The flaw area had a slightly higher hardness of R_C. 38-43 relative to the matrix hardness of R_C 34. This was most likely a Type II, Category 3 flaw (Fig. "Titanium flaw types"). The weak point was near the highly-stressed circumferential dove-tail groove.
The oxygen accumulation occured during the triple remelting process. A possible cause was a pronounced vacuum loss during the second remelting process. This occurrence could be reconstructed with the aid of the required documented process records (Fig. "Segregations influenced by melting temperature"). It is rare for such a process to occur with unallowable intensity. It occurs when the electrode shifts to a cooling water leak in the mold. However, the pressure increase was still within the limits that were tolerable at the time the raw material was produced (1972). The oxygen from the water, which dissociated at the high temperature, then diffuses into the melt bath (Fig. "Flaw causing titanium rotor fracture 2").
Typical characteristics of such a procedural deviation were attributed to the damaged part. This made it possible to identify other potentially threatened parts that were located near the damaged part in the casting block (also see Ill.15.2-22). In compliance with the demands of the responsible authorities, all suspect parts (21) were removed within 30 days.
This type of weak point with no cracking could evidently not be found through ultrasonic testing due to its unique location. The flaw, insofar as it spread to the surface, would have been detectable with the aid of macro-etching (blue etch anodizing = BEA). However, this process was not yet in use at the time.
Even in the finished part with potential cracking, ultrasonic testing during an overhaul or inspection was problematic. The flaws in the groove area were in a location that was most unfavorable for ultrasonic testing (blind spot). This is especially true of cracks that are located directly below the surface. During the last overhaul, ultrasonic testing showed anomalous findings, but they were categorized as allowable in accordance with regulations.
For this reason, additional eddy-current testing was recommended for the detection of cracks near the surface.

Figure "Flaw causing titanium rotor fracture 2" (Ref. 15.2-18): The optimization of the manufacturing process for aviation-suitable raw parts made from high-strength titanium alloys takes years. Today, remelting is usually repeated three times in a vacuum. The following describes typical process steps:

In step 1, pure titanium sponge is mixed with the powdered alloy components. This is pressed into large cuboids which are welded under shielding gas with strips into the first electrode with a diameter of 45cm (measurements of typical example; left diagram). An intermediate piece connects the electrode with the header.
This electrode is melted by an arc into a water-cooled mold, where it becomes a second electrode with a diameter of 75 cm.(see Ills. 15.1-11 and 15.1-12). Because material can melt from the intermediate piece during this process, it must also be made from rotor-quality titanium, as must the connecting straps.
For the second remelting, three of the ingots created in step 1, in this case with a diameter of 60 cm, are combined into an electrode (middle diagram). In the second remelting (step 2), unlike in step 1, no head material may be melted. A melting process similar to that in step 1 results in an ingot with a diameter of 75 cm.

The resulting ingot is inverted and creates the electrode for the third remelting (step 3). This creates the final product of the melting process, which is an ingot with 90 cm diameter.
This ingot is turned, heat-treated, and forged into billets (see Fig. "Terms of part production stages"). It is then split into lengths that are required for the specific raw part volume. Machining creates the geometry necessary for the forming process into part-specific blanks (e.g. drop-forging part).

Notes for Fig. "Flaw causing titanium rotor fracture": Originally, the blanks for rotors in stages 3-9 (middle right diagram) were made from a single pre-forged billet with a diameter of 40 cm. Then, billets with diameters of 30 and 32.5 cm were used. Later, two 25 cm billets were used, and these were eventually replaced by 20 cm billets. The two parts were prefabricated and then welded into a rotor. The major re-forming of the thin billets into a forging blank leads to a more favorable micro-structure, and reduces flaw sizes and changes their orientation. The trend towards ever smaller billet diameters was also intended to aid detection of hard a segregations. It created a larger surface that could be inspected through etching/BEA (blue etch anodizing).

Figure "Problems of titanium alloy rotor raw parts" (Ref. 15.2-30, Example 15.2-4): After 41,009 operating hours with 15,503 startup/shutdown cycles, an LCF fracture occurred in the fan disk (middle diagram), which was made of the titanium alloy TiAl6V4. 760 cycles before the damage, the engine was in for repair, in the course of which the disk was inspected with fluorescent penetrant. An ultrasonic inspection of the dovetail grooves in the annulus was also conducted. No cracks were reported by these inspections.
The crack that led to the fracture originated in a 1.4 x 0.3 mm flaw (left bottom detail) near the hub bore. The cyclical (LCF) crack growth between 14,000 and 16,000 load cycles led to the critical crack length of roughly 30 mm at the surface of the hub bore. In the crack initiation zone, there was an accumulation of hard alpha material with a high nitrogen content. In this zone, there were also microcracks and microporosity. These are typical characteristics of flaw type I (Fig. "Titanium flaw types"). In a “sister disk” from the same melting batch, ultrasonic inspection revealed anomalies in the disk membrane of the hub bore and the flange arm. They were the same flaw type as in the damaged disk. However, there was no LCF crack growth. In addition, macro-etching (blue etch anodizing) showed segregations that could be categorized as flaw type II with a high aluminum content.
The flaws with the increased nitrogen content were evidently brought in during the remelting of the ingot (Ill. 15.2-30).

Figure "Residual process stresses in titanium rotor disk" (Ref. 15.2-17, Example 15.2-5): Disk fractures occurred in a large first-generation fan engine. The material was a high-strength titanium alloy. Operating cracks (LCF) in several other disks that were subsequently inspected were causallly attributed to problems during raw part production. Because the fragments could not be recovered, the inspection results of the parts with operating cracks are treated as representative of the accident cases. This means that these cracks are assumed to correspond to the damage mechanism of the disk fractures.
Evidently, tension residual stresses from the forging process (bottom right diagram) overlayed with high operating loads in the affected disk area.
In addition, embrittling structural flaws with microcracks and microporosity (Fig. "Titanium flaw types") were present. They originated in the upper cast bar zone (Fig. "Flaw causing titanium rotor fracture 2") and remained in the pre-forged material (A-billets, bottom left diagram). The LCF crack probably originated in this type of flaw.

Figure "Pores causing cracks in titanium alloys" (Ref. 15.2-17): Pores in titanium forged parts, such as in this compressor disk, can be traced back to the casting process during semi-finished part production (Fig. "Titanium flaw types"). In the depicted case, a pore located near a bolt bore in the disk membrane led to cracking that spread to the surface. This crack was found with penetrant testing.
The pores are individual micropores (top left detail) or pore clusters. These pores promote cyclical crack growth in several different ways:

Pores can act as notches and increase local operating stresses considerably.

If the pores are connected with hydrogen absorption, it can result in embrittling effects with spontaneous cracking. This can even happen in new parts if the residual stresses are high enough.

If the pores are surrounded by hard brittle structure zones (“hard a), which are usually caused by oxygen or nitrogen absorption, it promotes brittle micro-cracking (Fig. "Problems of titanium alloy rotor raw parts") and increases the crack growth rate.

At the mouth of operating cracks, there is often a pore which is located in the center of a brittle, circular fracture plane (cleavage crack, quasi-cleavage crack). This fracture plane will not necessarily have crack growth lines.

Pore cracking has been observed for decades in fusion welds of titanium parts (electron beam, Ills. 16.2.1.3-30 and 16.2.1.3-31; shielding gas; Ref. 15.2-24). The cracking at the pore is largely determined by the residual stress. A hydrogen content greater than 200 ppm promotes crack initiation. An influence of oxygen has also been observed. Experts evidently agree that the brittle cracking at pores in titanium is causally related to the absorption of gases. The cracking is time-dependent, which indicates diffusion processes. It was thought that cracking had been observed during cold forming (trueing) of welds in thick-walled parts such as pressure vessels. However, this is contradicted by the expectation that these conditions would rather result in a fracture surface with typical ductile micro-characteristics (lappeting).
Today, it is evidently assumed that pore cracking in engine parts is related to dwell time fatigue (see Volume 3, Ills. 12.6.1-20, 15.2-18, and 15.2-19). This is a damage mechanism that requires a combination of static loads (constant rpm, residual stress) and cyclical loads (startup/shutdown cycles).

Figure "Melting segreations cause turbine disk failure" (Refs. 15.2-25 and 15.2-26, Example 15.2-6): After startup, a turbine disk made from a forged Ni alloy (Incoloy 901) fractured. It was the third low-pressure turbine stage (see engine diagram). The following inspection confirmed that an LCF crack had formed in the hub bore during operation. The crack grew radially outward to a blade groove (right diagram). A count of the crack growth lines indicated that there must already have been a crack present along the hub bore before the last overhaul. This calls into question the sufficient safety of the regulation-specified penetrant testing.
Similar cracks had already been discovered in two disks before the accident (Fig. "Residual process stresses in titanium rotor disk", Ref.17.2-17). However, these cases did not result in disk fractures. In order to bind the oxygen in the melt, oxidizing additives of cerium and lanthanum were used in all cases. These oxides are intended to float to the top of the melt bath (Fig. "White spots in Ni alloys"). They collect at the top of the ingot as dross, and are cut off before further processing. However, the oxides evidently led to a segregation in the highly-stressed disk hub zone. After this was discovered, the addition of cerium was stopped. Disks that were made from raw material produced with the same methods had to be removed and non-destructively tested. The test must be repeated every 5,900 hours.
In the case of the accident disk, subsequent investigation revealed that it was located below the dross in the ingot (bottom left diagram).

Figure "Segregation induced turbine disk failure": This engine damage occurred during the first test run of the engine, which was installed in a new aircraft. It is a single-shaft engine of an older type (middle diagram). A section of the first turbine stage disk broke out after a few minutes of total run time (bottom right diagram). The relatively small disk fragment with a blade pair broke through the turbine housing and escaped from the nacelle.
Subsequent inspection of the damaged disk revealed a large, dark oxidized flaw several square centimeters in size in the fracture zone. However, the flaw had not penetrated to the surface before the fracture. The flaw had a typical bow shape corresponding to the grain direction (Fig. "Flaws in forged rotor disks"). The surface of the residual force fracture was very small, because the flaw was located very close to the surface on both sides of the annulus.
A laboratory inspection revealed that the flaw was a large carbon nitride segregation. Experience has shown that this type of impurity can result in the disk material in question (hardenable iron-based alloy A286) during the casting process (Fig. "White spots in Ni alloys").
Further research at the raw part manufacturer revealed that the disk was made from material located directly below the removed dross. This was exceptionally long in accordance with the applicable procedural controls. This indicates an increased flaw risk. Evidently, the amount of removed material was insufficient.
Macro-etching tests of the neighboring disks from the remelt ingot also revealed segregations. The greater the distance from the ingot head, the less pronounced the flaws were, and the farther inside they were located (bottom right diagram, Fig. "Segregation distributions in billets"). This observation can be explained with the typical flaw distribution in the ingot head (Fig. "Terms of part production stages") and the subsequent forging process.
The prescribed penetrant test was not able to detect the internal flaw. Unfortunately, the ultrasonic testing process stipulated in the specifications was done in such a way that the the echo from the curved surface was not reflected to the receiver, which was positioned separately in this case (Ills. 17.3.1-5 and 17.3.1-9).

Figure "Dirty whie spots in Ni-base forging" (Ref. 15.2-28): This is a spinning disk made of a high-strength Ni-based forged alloy (Waspaloy). After vacuum induction melting (VIM), the material was remelted two times (DEVR), extruded, and heat-treated:

• Annealed for 4 hours at 995°C - 1035°C and then quenched in oil.
• Annealed for 4 hours at 850°C, then cooled in air.
• Annealed for 16 hours at 760°C, then cooled in air.

The centrifugal test was done at 500°C.
The test was aborted after 3,734 cycles. A large LCF crack (3.81 x 15.24 mm) had formed from the edge of a bolt bore (left diagram). A laboratory inspection revealed that the LCF crack (bottom right diagram) spread from a 1.73 x 1.85 mm non-metallic inclusion (aluminum oxide, top right diagram).
The text does not explain how this impurity entered the material during the casting or remelting processes. However, because this is evidently a rather massive Al₂O₃ particle, it may have come from a ceramic filter for the melt or from lining material during melting of the ingot.

Figure "Material separations in forgings": A special flaw type in rolled and forged parts is two-dimensional separations that tend to run parallel to the surface. These flaws were found more frequently until the late 1960s. Today, they are evidently extremely rare. This is most likely related to improved non-destructive testing and more advanced materials technology. The separations can occur through very different mechanisms. Their shape is determined by the forming process from the cast ingot to the semi-finished/raw part. Typical flaws of this type are:

Laminations: These occur in rolled materials such as profiles and sheets. They are gas bubbles from the casting process that have been rolled flat (bottom right diagram). During heat treatment at high temperatures, especially in a vacuum, a bulge can form at the flaw. The causes are the gas pressure in the bubble and/or compressive stress caused by the more rapid expansion of the unconnected surface layer (top left diagram).
A similar effect can occur when several sheets are fused into a larger section through rolling, and local bonding problems occur. The findings in the top right detail correspond to this process. The material is a high-alloy hardenable austenitic steel (A286).

Forging laps: Occur during drop forging when material is sheared off and “smeared” by the forging motion. A typical example is precision-forged blade profiles (bottom right diagram). In this case, material from the blade root platform is transferred to the top of the blade by the forging motion or top half of the die. The depicted case concerns a compressor rotor blade made from a type of 13% Cr steel.

Separation through segregations: Larger segregations can be broken open through forging deformation and/or thermal stress. The impurity prevents the separation from resealing during subsequent forming processes. The result is internal separations that are formed into the plane of the grain direction by the forging process (Ills. 15.1-13 and 15.2-11).

Separations at unforged cracks: If cracks occur ahead of the last forging process (Fig. "Flaws in forged rotor disks") and do not reseal, it is possible that laminar separations parallel to the surface will result. These cracks can be the result of overstressing of the material during forging. Excessive shear loads can break open the surface. Oxidation of the separation makes it impossible to fuse it shut during subsequent forging cycles. (Thermal) cracks can also occur inside forged parts. The forged part reaches a high temperature due to the applied forming energy. If this results in overtemperature and damage-promoting softening of the grain boundaries, the material can break open under the strain of the forming and/or thermal stress. However, these oxide-free separation surfaces will usually re-fuse in a subsequent forging process.
Cracks can also be caused by overly high thermal stress. If they come into contact with oxygen, resealing through subsequent forging is no longer possible.

Figure "Stress corrosion cracking in brazed steel parts": In older engine types, Cr steels were used in the compressor blading. These are heat-treated steels. The top right diagram shows a “constructed” guide vane. The blade, made from a rolled material, is soldered onto the root box, which is made of sheet metal. During operation, cracks occurred due to stress corrosion cracking (SCC, Volume 1, Chapter 5.4.2) in the root boxes (top right diagram). The cause for the susceptibility to cracking was evidently that the annealing temperature was too low during the production process (middle right diagram). The results were excessive hardness and sensitivity to SCC (bottom right diagram). At the same time, the low annealing temperature resulted in dangerously high tension residual stresses, due to the effective high yield limit (bottom left diagram).

Figure "Production related crackung of blades" (Ref. 15.2-29): In two cases, uncontained blade fractures (right diagram) occurred in fan engines of an older type. The affected fan blades were made of a high-strength titanium alloy. Laboratory tests revealed that the fractures were HCF dynamic fatigue fractures. The fractures originated in a discolored zone below the clapper. The crack initiation zone was characterized by intercrystalline cracking and transcrystalline fracture planes. These damage symptoms are typical in this high-strength titanium alloy for stress corrosion cracking (SCC) under hot salt. Because this damage mechanism requires minimum temperatures of around 500°C (Volume 1, Ill. 5.4.2.1-8), the crack could not have occurred during operation (maximum temperature 116 °C). It was revealed that hand sweat from the blade production process was evidently the cause. The sweat was on a section of the blade that had dangerously high undetected tension residual stress from the forging process. The final heat treatment at 538 °C then caused the SCC. The cracks were evidently not detected by the non-destructive testing methods used in this case.
The left diagram shows a compressor rotor blade from an older engine design, made from a type of 13% Cr steel. The undeformed, darkly colored crack through the root shaft can probably be traced back to the forging process. It is surprising that such a flaw was not detected despite the sensitive magnetic crack detection done on the new part.